U.S. patent application number 14/211067 was filed with the patent office on 2014-09-25 for cu-ti based copper alloy sheet material and method for producing the same, and electric current carrying component.
This patent application is currently assigned to DOWA METAL TECH CO., LTD.. The applicant listed for this patent is DOWA METAL TECH CO., LTD.. Invention is credited to Weilin GAO, Toshiya Kamada, Takashi Kimura, Fumiaki Sasaki, Akira Sugawara, Motohiko Suzuki.
Application Number | 20140283963 14/211067 |
Document ID | / |
Family ID | 50423951 |
Filed Date | 2014-09-25 |
United States Patent
Application |
20140283963 |
Kind Code |
A1 |
GAO; Weilin ; et
al. |
September 25, 2014 |
Cu-Ti BASED COPPER ALLOY SHEET MATERIAL AND METHOD FOR PRODUCING
THE SAME, AND ELECTRIC CURRENT CARRYING COMPONENT
Abstract
A Cu--Ti based copper alloy sheet material contains, in mass %,
from 2.0 to 5.0% of Ti, from 0 to 1.5% Ni, from 0 to 1.0% Co, from
0 to 0.5% Fe, from 0 to 1.2% Sn, from 0 to 2.0% Zn, from 0 to 1.0%
Mg, from 0 to 1.0% Zr, from 0 to 1.0% Al, from 0 to 1.0% Si, from 0
to 0.1% P, from 0 to 0.05% B, from 0 to 1.0% Cr, from 0 to 1.0% Mn,
and from 0 to 1.0% V, the balance substantially being Cu. The sheet
material has a metallic texture wherein in a cross section
perpendicular to a sheet thickness direction, a maximum width of a
grain boundary reaction type precipitate is not more than 500 nm,
and a density of a granular precipitate having a diameter of 100 nm
or more is not more than 10.sup.5 number/mm.sup.2.
Inventors: |
GAO; Weilin; (Tokyo, JP)
; Suzuki; Motohiko; (Tokyo, JP) ; Kamada;
Toshiya; (Tokyo, JP) ; Kimura; Takashi;
(Tokyo, JP) ; Sasaki; Fumiaki; (Tokyo, JP)
; Sugawara; Akira; (Tokyo, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
DOWA METAL TECH CO., LTD. |
Tokyo |
|
JP |
|
|
Assignee: |
DOWA METAL TECH CO., LTD.
Tokyo
JP
|
Family ID: |
50423951 |
Appl. No.: |
14/211067 |
Filed: |
March 14, 2014 |
Current U.S.
Class: |
148/685 ;
148/411; 148/412; 148/413; 148/414 |
Current CPC
Class: |
C22F 1/08 20130101; H01B
1/026 20130101; C22C 9/00 20130101 |
Class at
Publication: |
148/685 ;
148/411; 148/412; 148/413; 148/414 |
International
Class: |
H01B 1/02 20060101
H01B001/02; C22C 9/00 20060101 C22C009/00; C22F 1/08 20060101
C22F001/08 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 25, 2013 |
JP |
2013-061498 |
Claims
1. A copper alloy sheet material which comprises from 2.0 to 5.0%
of Ti, from 0 to 1.5% of Ni, from 0 to 1.0% of Co, from 0 to 0.5%
of Fe, from 0 to 1.2% of Sn, from 0 to 2.0% of Zn, from 0 to 1.0%
of Mg, from 0 to 1.0% of Zr, from 0 to 1.0% of Al, from 0 to 1.0%
of Si, from 0 to 0.1% of P, from 0 to 0.05% of B, from 0 to 1.0% of
Cr, from 0 to 1.0% of Mn, and from 0 to 1.0% of V in terms of % by
mass, with a total content of Sn, Zn, Mg, Zr, Al, Si, P, B, Cr, Mn,
and V being not more than 3.0% and the balance being Cu and
inevitable impurities, wherein the copper alloy sheet material has
a metallic texture in which in a cross section thereof
perpendicular to the sheet thickness direction, a maximum width of
a precipitate of grain boundary reaction type is not more than 500
nm, and a density of a granular precipitate having a diameter of
100 nm or more is not more than 10.sup.5 number/mm.sup.2.
2. The copper alloy sheet material according to claim 1, wherein
the metallic texture further has an average crystal grain diameter
of from 5 to 25 .mu.m in a cross section thereof perpendicular to
the sheet thickness direction.
3. The copper alloy sheet material according to claim 1, wherein
the copper alloy sheet material has an electrical conductivity of
15% IACS or more.
4. The copper alloy sheet material according to claim 1, wherein
when the rolling direction of the sheet is defined as LD, and the
direction rectangular to the rolling direction and the sheet
thickness direction is defined as TD, the copper alloy sheet
material has a 0.2% offset yield strength in LD of 850 MPa or more
and has bending workability such that in the 90.degree. W-bending
test in conformity with JIS H3130, a value of R/t ratio of a
minimum bending radius R to a sheet thickness t at which cracking
does not occur is not more than 2.0 in both LD and TD.
5. The copper alloy sheet material according to claim 1, wherein
the copper alloy sheet material has fatigue resistance such that in
the fatigue test in conformity with JIS Z2273, a test piece in
which the rolling direction of the sheet is the longitudinal
direction has a fatigue life at a maximum load stress of 700 MPa on
the test piece surface (the number of repeated vibrations until
rupture of the test piece occurs) is 500,000 times or more.
6. A method for producing a copper alloy sheet material according
to claim 1, comprising a step of subjecting a sheet material having
been subjected to hot rolling and cold rolling at a rolling ratio
of 90% or more, to a heat treatment with a heat pattern including a
solution treatment at from 750 to 950.degree. C., holding at a
temperature ranging from 550 to 730.degree. C. in a cooling process
after the solution treatment for from 10 to 120 seconds, and then
quenching to at least 200.degree. C. at an average cooling rate of
20.degree. C./sec or more; and a step of subjecting the sheet
material after the heat treatment successively to intermediate cold
rolling at a rolling ratio of from 0 to 50%, an aging treatment at
from 300 to 430.degree. C., and finish cold rolling at a rolling
ratio of from 0 to 30%.
7. A method for producing a copper alloy sheet material according
to claim 1, comprising a step of subjecting a sheet material having
been subjected to hot rolling and cold rolling at a rolling ratio
of 90% or more, to a heat treatment with a heat pattern including a
solution treatment at from 750 to 950.degree. C., then quenching to
at least 200.degree. C. at an average cooling rate of 20.degree.
C./sec or more, thereafter increasing the temperature and holding
at a temperature ranging from 550 to 730.degree. C. for from 10 to
120 seconds, and then quenching to at least 200.degree. C. at an
average cooling rate of 20.degree. C./sec or more; and a step of
subjecting the sheet material after the heat treatment successively
to intermediate cold rolling at a rolling ratio of from 0 to 50%,
an aging treatment at from 300 to 430.degree. C., and finish cold
rolling at a rolling ratio of from 0 to 30%.
8. The method for producing a copper alloy sheet material according
to claim 6, wherein the method includes controlling a rolling ratio
of the finish cold rolling to from 5 to 30% and then applying
low-temperature annealing at from 150 to 430.degree. C.
9. The method for producing a copper alloy sheet material according
to claim 6, wherein the method includes adjusting a heating time
and an in-furnace time in the solution treatment such that an
average crystal grain diameter in a cross section perpendicular to
the sheet thickness direction after the final cold rolling is from
5 to 25 .mu.m.
10. An electric current carrying component using the copper alloy
sheet material according to claim 1 for a material.
Description
TECHNICAL FIELD
[0001] The present invention relates to a Cu--Ti BASED copper alloy
sheet material suitable for electric current carrying components
such as connectors, lead frames, relays, and switches, and in
particular, the present invention relates to a sheet material
having conspicuously improved fatigue resistance and a method for
producing the same. In addition, the present invention relates to
an electric current carrying component using the copper alloy sheet
material for a material.
BACKGROUND ART
[0002] Materials which are used for electric current carrying
components constituting electrical or electronic components such as
connectors, lead frames, relays, and switches are required to have
high "strength" capable of withstanding a stress which is given at
the time of assembling or operation of an electrical or electronic
appliance. In addition, materials used for electrical or electronic
components are required to have excellent "bending workability"
because said components are generally formed by bending.
Furthermore, in order to ensure contact reliability between
electrical or electronic components, durability against a
phenomenon in which a contact pressure decreases with time (stress
relaxation), namely excellent "stress relaxation resistance" is
required, too. The stress relaxation as referred to herein is a
kind of creep phenomenon in which even if the contact pressure of a
spring part of an electric current carrying component constituting
an electrical or electronic component is kept in a fixed state at
ordinary temperature, it decreases with time under an environment
of relatively high temperatures (for example, from 100 to
200.degree. C.). That is, the stress relaxation means a phenomenon
in which in a state where a stress is given to a metal material,
dislocation moves due to self-diffusion of atoms constituting the
matrix or diffusion of a solute atom to cause plastic deformation,
whereby the given stress is relieved. When used in an environment
where an increase of the component temperature as in automobile
connectors is supposed, the "stress relaxation resistance" is
particularly important.
[0003] In the light of the above, materials which are used for
electrical or electronic components are required to have excellent
"strength", "bending workability", and "stress relaxation
resistance". Meanwhile, electric current carrying components having
a movable portion, such as relays and switches, are also required
to have excellent "fatigue resistance" in terms of durability
capable of withstanding a repeated stress load. But, in general,
the "fatigue resistance" or "bending workability" is in a trade-off
relationship with the "strength", and in copper alloy sheet
materials, it is not easy to enhance the "fatigue resistance" or
"bending workability" at the same time while contemplating to
achieve high strength.
[0004] Among copper alloys, a Cu--Ti based copper alloy has high
strength just below a Cu--Be based copper alloy and has stress
relaxation resistance superior to the Cu--Be copper alloy. In
addition, the Cu--Ti based copper alloy is more advantageous than
the Cu--Be copper alloy from the standpoints of cost and
environmental load. For that reason, the Cu--Ti based copper alloy
(for example, C1990 which is a Cu-3.2% by mass Ti alloy) is used
for connector materials or the like as an alternate material of the
Cu--Be based copper alloy. But, the Cu--Ti based copper alloy is
generally inferior in the "fatigue resistance" and "bending
workability" to the Cu--Be copper alloy having equal strength.
CITATION LIST
Patent Literatures
[0005] Patent Literature 1: JP-A-2012-87343 ("JP-A" means
unexamined published Japanese patent application) [0006] Patent
Literature 2: JP-A-2012-97308
SUMMARY OF INVENTION
Problems to be Solved by the Invention
[0007] As is well known, the Cu--Ti based copper alloy is an alloy
capable of enhancing the strength utilizing a modulated structure
(spinodal structure) of Ti. The modulated structure is a structure
which is formed while keeping complete consistency with a mother
phase due to a continuous fluctuation in the concentration of a Ti
solute atom. Though the material is conspicuously hardened by the
modulated structure, a loss of the fatigue resistance or bending
workability to be caused due to this matter is relatively
small.
[0008] Meanwhile, Ti in the Cu--Ti based copper alloy mother phase
forms an intermetallic compound (.beta. phase) with Cu to
precipitate as a second phase grain in the crystal grain boundary
or the grain. In this specification, a granular precipitate
including an intermetallic compound of this kind is generically
named "granular precipitate". The greater part of the granular
precipitate observed in the Cu--Ti based copper alloy is a grain of
the above-described .beta. phase. In addition, when Ti in the
mother phase reacts with Cu in the crystal grain boundary, a
striped intermetallic compound precipitates from the grain boundary
and grows. The intermetallic compound phase of this kind is named
"precipitate of grain boundary reaction type".
[0009] The granular precipitate is small in its own hardening
action, so that when a large amount of the granular precipitate
precipitates, it brings a decrease of the concentration of a solute
Ti atom constituting a modulated structure, thereby becoming a
factor to hinder the enhancement of strength. In addition, the
precipitate of grain boundary reaction type is a weak portion and
easily becomes a starting point of fatigue fracture. Patent
Literature 2 discloses a technology for improving the strength,
electrical conductivity, and bending workability by increasing an
existent ratio of a precipitate of grain boundary reaction type
which is occupied in a precipitation phase in a Cu--Ti based copper
alloy. It is said that coarsening of a stable phase (granular
precipitate) is suppressed due to the formation of a precipitate of
grain boundary reaction type, and as a result, a 0.2% offset yield
strength of 850 MPa or more can be realized while suppressing a
decrease of the bending workability. But, according to
investigations made by the present inventors, the precipitate of
grain boundary reaction type is originally a weak portion, and the
precipitate of grain boundary reaction type itself becomes a factor
to decrease the strength or bending workability. In particular, in
order to improve the fatigue resistance, it is necessary to
suppress the formation of a precipitate of grain boundary reaction
type.
[0010] In the case of a Cu--Be based copper alloy, by adding Co or
Ni, such an additive element segregates in the grain boundary,
thereby making it possible to suppress the precipitation of grain
boundary reaction type. However, in the Cu--Ti based copper alloy,
in view of the fact that Ti is a very active element, the additive
element is easily consumed through the formation of a compound with
Ti, so that an effect for suppressing the precipitation of grain
boundary reaction type utilizing the segregation in the grain
boundary is small. In addition, a primary strengthening mechanism
of the Cu--Ti based copper alloy is one derived from a modulated
structure (spinodal structure) of solid-solved Ti, and therefore,
the addition of a large amount of a third element decreases the
amount of solid-solved Ti to offset the merits of the Cu--Ti based
copper alloy each other.
[0011] The precipitate of grain boundary reaction type of the
Cu--Ti based copper alloy is formed chiefly in an aging treatment
process. It is the present state that a technology for effectively
suppressing the formation of a precipitate of grain boundary
reaction type has not been established yet, and it may be
considered that it is difficult to enhance the fatigue resistance
of the Cu--Ti based copper alloy. The present invention is to
provide a Cu--Ti based copper alloy sheet material having improved
"fatigue resistance" while keeping the "strength", "bending
workability", and "stress relaxation resistance" good.
Means for Solving the Problems
[0012] The aging treatment temperature for bringing out a maximum
strength of the Cu--Ti based copper alloy is generally from about
450 to 500.degree. C. But, the precipitation of grain boundary
reaction type is simultaneously caused in this temperature region.
As a result of detailed investigations made by the present
inventors, it has been found that by performing a heat treatment in
a temperature region of from 550 to 730.degree. C. after a solution
treatment, a precursory texture state of the modulated structure is
obtained; and that in those having such a texture state, the aging
treatment temperature at which a maximum strength is obtained
shifts towards the low-temperature side. Specifically, it becomes
possible to achieve the aging treatment at low temperatures as from
300 to 430.degree. C. In that temperature region, the formation of
a precipitate of grain boundary reaction type can be effectively
suppressed. The present invention has been accomplished on the
basis of such knowledge.
[0013] Specifically, the above-described object is achieved by a
copper alloy sheet material which comprises from 2.0 to 5.0% of Ti,
from 0 to 1.5% of Ni, from 0 to 1.0% of Co, from 0 to 0.5% of Fe,
from 0 to 1.2% of Sn, from 0 to 2.0% of Zn, from 0 to 1.0% of Mg,
from 0 to 1.0% of Zr, from 0 to 1.0% of Al, from 0 to 1.0% of Si,
from 0 to 0.1% of P, from 0 to 0.05% of B, from 0 to 1.0% of Cr,
from 0 to 1.0% of Mn, and from 0 to 1.0% of V in terms of % by
mass, with a total content of Sn, Zn, Mg, Zr, Al, Si, P, B, Cr, Mn,
and V being not more than 3.0% and the balance being Cu and
inevitable impurities, wherein the copper alloy sheet material has
a metallic texture in which in a cross section thereof
perpendicular to the sheet thickness direction, a maximum width of
a precipitate of grain boundary reaction type is not more than 500
nm, and a density of a granular precipitate having a diameter of
100 nm or more is not more than 10.sup.5 number/mm.sup.2. A copper
alloy sheet material having a metallic texture in which in a cross
section thereof perpendicular to the sheet thickness direction, an
average crystal grain diameter is from 5 to 25 .mu.m is more
suitable for the subject. An electrical conductivity of 15% IACS or
more can be ensured. The maximum width of the precipitate of grain
boundary reaction type as referred to herein means a maximum value
of a length of the precipitate of grain boundary reaction type in
the rectangular direction to a crystal grain boundary on which the
precipitate of grain boundary reaction type is formed, the length
being measured at a position above the crystal grain boundary in
the observation of metallic texture. The "diameter" of the granular
precipitate means a major axis of the grain in the observation of
metallic texture.
[0014] In the above-described copper alloy sheet material, when the
rolling direction of the sheet is defined as LD, and the direction
rectangular to the rolling direction and the sheet thickness
direction is defined as TD, it is possible to realize a copper
alloy sheet material having a 0.2% offset yield strength in LD of
850 MPa or more and having bending workability such that in the
90.degree. W-bending test in conformity with JIS H3130, a value of
R/t ratio of a minimum bending radius R to a sheet thickness t at
which cracking does not occur is not more than 2.0 in both LD and
TD. In addition, with respect to the fatigue properties, it is
possible to provide a copper alloy sheet material having excellent
fatigue resistance such that in the fatigue test in conformity with
JIS 22273, in a test piece in which the rolling direction of the
sheet is the longitudinal direction, a fatigue life at a maximum
load stress of 700 MPa on the test piece surface (the number of
repeated vibrations until rupture of the test piece occurs) is
500,000 times or more. The above-described copper alloy sheet
material is extremely useful as a material for working into an
electric current carrying component. Though the sheet thickness of
the above-described copper alloy sheet material can be made to, for
example, from 0.05 to 1.0 mm, in order to respond to thin-wall
processing of an electric current carrying component, it is
preferable to make the sheet thickness of the copper alloy sheet
material to, for example, from 0.05 to 0.35 mm.
[0015] The above-described copper alloy sheet material can be
obtained by a production method including
[0016] a step of subjecting a sheet material having been subjected
to hot rolling and cold rolling at a rolling ratio of 90% or more,
to a heat treatment with a heat pattern including a solution
treatment at from 750 to 950.degree. C., holding at a temperature
ranging from 550 to 730.degree. C. in a cooling process after the
solution treatment for from 10 to 120 seconds, and then rapidly
cooling to at least 200.degree. C. at an average cooling rate of
20.degree. C./sec or more; and
[0017] a step of subjecting the sheet material after the heat
treatment successively to intermediate cold rolling at a rolling
ratio of from 0 to 50%, an aging treatment at from 300 to
430.degree. C., and finish cold rolling at a rolling ratio of from
0 to 30%.
[0018] In addition, after carrying out the solution treatment by
means of a usual step, a step of reheating at a temperature ranging
from 550 to 730.degree. C. can also be adopted as a pretreatment of
the aging treatment. In that case, there can be applied a
production method including
[0019] a step of subjecting a sheet material having been subjected
to hot rolling and cold rolling at a rolling ratio of 90% or more,
to a heat treatment with a heat pattern including a solution
treatment at from 750 to 950.degree. C., then quenching to at least
200.degree. C. at an average cooling rate of 20.degree. C./sec or
more, thereafter increasing the temperature and holding at a
temperature ranging from 550 to 730.degree. C. for from 10 to 120
seconds, and then quenching to at least 200.degree. C. at an
average cooling rate of 20.degree. C./sec or more; and
[0020] a step of subjecting the sheet material after the heat
treatment successively to intermediate cold rolling at a rolling
ratio e of from 0 to 50%, an aging treatment at from 300 to
430.degree. C., and finish cold rolling at a rolling ratio of from
0 to 30%.
[0021] The "rolling ratio of 0%" as referred to herein means that
the rolling is not carried out. That is, the intermediate cold
rolling or finish cold rolling can be omitted. In the case of
carrying out the finish cold rolling, it is preferable to adopt a
step of controlling its rolling ratio to from 5 to 30% and then
applying low-temperature annealing at from 150 to 430.degree. C. In
addition, it is preferable to adjust a heating time and an
in-furnace time in the solution treatment such that an average
crystal grain diameter in a cross section perpendicular to the
sheet thickness direction after the final cold rolling is from 5 to
25 .mu.m.
Advantages of the Invention
[0022] According to the present invention, it has become possible
to provide a Cu--Ti based copper alloy sheet material which is
excellent in strength, bending workability, and stress relaxation
resistance and is also excellent in fatigue resistance. The present
invention is useful for needs of downsizing and thin-wall
processing of electrical or electronic components, which will be
expected to be developed more and more in the future.
BRIEF DESCRIPTION OF THE DRAWINGS
[0023] FIG. 1 is an SEM photograph of metallic texture of a general
Cu--Ti based copper alloy.
[0024] FIG. 2 is an SEM photograph of metallic texture of
Comparative Example No. 21 produced in usual steps.
[0025] FIG. 3 is an SEM photograph of metallic texture of Example
No. 1 according to the present invention.
MODES FOR CARRYING OUT THE INVENTION
<<Alloy Composition>>
[0026] In the present invention, a Cu--Ti based copper alloy in
which a binary basic component of Cu--Ti is blended with Ni, Co,
Fe, and other alloying elements, if desired is adopted. The term
"%" regarding the alloy composition hereunder means "% by mass"
unless otherwise indicated.
[0027] Ti is an element having a high age hardening action in a Cu
matrix and contributes to an increase of the strength and an
enhancement of the stress relaxation resistance. In order to
sufficiently bring out these actions, it is advantageous to ensure
the Ti content of preferably 2.0% or more, and more preferably 2.5%
or more. On the other hand, when the Ti content is in excess,
cracking easily occurs during a hot working or cold working
process, and a lowering of productivity is easily brought. In
addition, the temperature region in which the solution treatment
can be achieved becomes narrow, so that it becomes difficult to
bring out good properties. As a result of various investigations,
it is necessary to control the Ti content to not more than 5.0%.
The Ti content is adjusted within the range of preferably not more
than 4.0%, and more preferably not more than 3.5%.
[0028] Each of Ni, Co, and Fe is an element which contributes to an
enhancement of the strength upon formation of an intermetallic
compound with Ti, and at least one member of these elements can be
added, if desired. In particular, in the solution treatment of the
Cu--Ti based copper alloy, since such an intermetallic compound
suppresses coarsening of the crystal grain, it is possible to carry
out the solution treatment in a higher temperature region, and such
is advantageous in sufficiently achieving solid-solution of Ti. As
to the content in the case of adding at least one member of these
elements, it is more effective to contain 0.05% or more of Ni,
0.05% or more of Co, and 0.05% or more of Fe, respectively, and it
is still more effective to contain 0.1% or more of Ni, 0.1% or more
of Co, and 0.1% or more of Fe, respectively. However, when each of
Fe, Co, and Ni is contained in excess, the amount of Ti which is
consumed by the formation of the resulting intermetallic compound
becomes large, so that the amount of solid-solved Ti becomes small
inevitably. In that case, conversely, a lowering of the strength is
easily brought. In consequence, in the case of adding at least one
member of Ni, Co, and Fe, the contents of Ni, Co, and Fe are
controlled to the ranges of not more than 1.5%, not more than 1.0%,
and not more than 0.5%, respectively. The contents of Ni, Co, and
Fe may also be controlled to the ranges of not more than 0.25%, not
more than 0.25%, and not more than 0.25%, respectively.
[0029] Sn has an action to strengthen solid solution and an action
to enhance stress relaxation resistance. It is more effective to
ensure the Sn content of 0.1% or more. However, when the Sn content
exceeds 1.0%, castability and electrical conductivity are
conspicuously lowered. For that reason, in the case of containing
Sn, it is necessary to control the Sn content to not more than
1.0%. The Sn content may also be controlled to the range of not
more than 0.5%, or not more than 0.25%.
[0030] Zn has actions to enhance soldering properties and strength,
and besides, it also has an action to improve castability.
Furthermore, in the case of containing Zn, there is brought such an
advantage that an inexpensive brass scrap can be used. However, an
excess of the Zn content easily becomes a factor to cause a
lowering of electrical conductivity or stress corrosion cracking
resistance. For that reason, in the case of containing Zn, it is
necessary to control the Zn content to not more than 2.0%, and the
Zn content may also be controlled to the range of not more than
1.0%, or not more than 0.5%. In order to sufficiently obtain the
above-described actions, it is desirable to ensure the Zn content
of 0.1% or more, and in particular, it is more effective to control
the Zn content to 0.3% or more.
[0031] Mg has an action to enhance stress relaxation resistance and
a desulfurizing action. In order to sufficiently exhibit these
actions, it is preferable to ensure the Mg content of 0.01% or
more, and it is more effective to control the Mg content to 0.05%
or more. However, Mg is an element which is easily oxidized, and
when the Mg content exceeds 1.0%, castability is conspicuously
lowered. For that reason, in the case of containing Mg, it is
necessary to control the Mg content to not more than 1.0%, and it
is more preferable to adjust the Mg content within the range of not
more than 0.5%. In general, the Mg content may be controlled to not
more than 0.1%.
[0032] As other elements, it is possible to contain at least one
member of Zr of not more than 1.0%, Al of not more than 1.0%, Si of
not more than 1.0%, P of not more than 0.1%, B of not more than
0.05%, Cr of not more than 1.0%, Mn of not more than 1.0%, and V of
not more than 1.0%. For example, each of Zr and Al is able to form
an intermetallic compound with Ti, and Si is able to form a
precipitate with Ti. Each of Cr, Zr, Mn, and V easily forms a high
melting-point compound with S, Pb, or the like which exists as an
inevitable impurity. In addition, each of Cr, B, P, and Zr has a
refining effect of the cast texture and may contribute to an
improvement of hot workability. In the case of containing at least
one of Zr, Al, Si, P, B, Cr, Mn, and V, in order to sufficiently
obtain the actions of the respective elements, it is effective to
contain such an element in an amount of 0.01% or more in total.
[0033] However, what a large amount of at least one member of Zr,
Al, Si, P, B, Cr, Mn, and V is contained adversely affects the hot
or cold workability and is disadvantageous from the standpoint of
cost. In consequence, it is desirable to control a total content of
Sn, Zn, Mg, Zr, Al, Si, P, B, Cr, Mn, and V to not more than 3.0%.
The total content can be controlled to the range of not more than
2.0% or not more than 1.0%, and it may be controlled to the range
of not more than 0.5%. As for a more rational upper limit taking
the economy into account, for example, it can be controlled to not
more than 0.2% for Zr, not more than 0.15% for Al, not more than
0.2% for Si, not more than 0.05% for P, not more than 0.03% for B,
not more than 0.2% for Cr, not more than 0.1% for Mn, and not more
than 0.2% for V, respectively.
<<Metallic Texture>>
[0034] An SEM photograph of metallic texture of a general Cu--Ti
based copper alloy is illustrated in FIG. 1. A "granular
precipitate" of a type shown by a symbol A, and a "precipitate of
grain boundary reaction type" of a type shown by a symbol B are
observed. However, a strengthening mechanism of the Cu--Ti based
copper alloy is one mainly derived from a modulated structure
(spinodal structure). Different from the precipitate, the modulated
structure itself is not observed by an optical microscope or
SEM.
[Granular Precipitate]
[0035] As for the granular precipitate observed in a mother phase
(matrix) of the Cu--Ti based copper alloy, though intermetallic
compounds such as Ni--Ti, Co--Ti, and Fe--Ti intermetallic
compounds may be existent depending upon the kind of the alloying
element to be added, a .beta. phase that is a Cu--Ti intermetallic
compound occupies the majority in quantity. In the case where the
grain diameter of the granular precipitate is small as, for
example, from several nm to several tens of nm, not only the
hardening action effectively reveals, but a loss of ductility is
small. On the other hand, in a granular precipitate having a
diameter of 100 nm or more, nevertheless the hardening action is
small, a loss of ductility is large. In addition, when a large
amount of such a coarse granular precipitate is formed, the
concentration of a Ti solute atom in the modulated structure
decreases, and a lowering of the strength is caused. As a result of
various investigations, it is necessary to control a density of a
granular precipitate having a diameter of 100 nm or more to not
more than 10.sup.5 number/mm.sup.2, and the density of a granular
precipitate having a diameter of 100 nm or more is more preferably
not more than 5.times.10.sup.4 number/mm.sup.2.
[Precipitate of Grain Boundary Reaction Type]
[0036] According to investigations made by the present inventors,
the precipitate of grain boundary reaction type is a very weak
portion and becomes a factor to bring a lowering of the strength or
a lowering of the stress relaxation resistance. In addition, the
precipitate of grain boundary reaction type becomes a starting
point of fatigue fracture or bending cracking. In particular, in
order to improve the stress relaxation resistance, it has been
noted that it is extremely effective to strictly control the
formation amount of the precipitate of grain boundary reaction
type. As a result of detailed studies, when in a cross section
perpendicular to the sheet thickness direction, a maximum width of
the precipitate of grain boundary reaction type is not more than
500 nm, it is possible to stably realize excellent fatigue
resistance such that a fatigue life at a maximum load stress of 700
MPa in the fatigue test in conformity with JIS Z2273 is 500,000
times or more. The maximum width of the precipitate of grain
boundary reaction type is more preferably not more than 300 nm.
[0037] It is meant by the terms "in a cross section perpendicular
to the sheet thickness direction, a maximum width of the
precipitate of grain boundary reaction type is not more than X nm"
that in the cross section perpendicular to the sheet thickness
direction, namely in the observed surface of metallic texture
prepared by polishing the sheet surface, in the case where a length
of the precipitate of grain boundary reaction type is measured in
the rectangular direction to the crystal grain boundary in a
crystal grain boundary portion where the precipitate of grain
boundary reaction type is formed, a maximum value of the foregoing
length does not exceed X nm. The texture state in which the maximum
width of the precipitate of grain boundary reaction type is not
more than 500 nm or not more than 300 nm can be realized by
production steps including a "precursory treatment" as described
later.
[Average Crystal Grain Diameter]
[0038] What the average crystal grain diameter is smaller is
advantageous for enhancing the bending workability. In the case
where great importance is attached to the bending workability, the
average crystal grain diameter of a final product sheet material is
desirably not more than 25 .mu.m. The average crystal grain
diameter is adjusted to preferably not more than 20 .mu.m, and more
preferably not more than 15 .mu.m. On the other hand, when the
average crystal grain diameter is too small, the stress relaxation
resistance is easily lowered. As a result of various
investigations, in order to ensure a stress relaxation resistance
level on which high evaluation is obtainable in an application of
an onboard connector, the average crystal grain diameter of the
final product sheet material is desirably 5 .mu.m or more, and more
desirably 8 .mu.m or more. Control of the average crystal grain
diameter can be mainly carried out by a solution treatment. The
average crystal grain diameter can be determined by measuring the
grain diameter of 100 or more crystal grains in a visual field of
300 .mu.m.times.300 .mu.m or more in the observation of metallic
texture of the cross section perpendicular to the sheet thickness
direction by the cutting method of JIS H0501.
<<Properties>>
[Electrical Conductivity]
[0039] When needs of weight-saving and thin-wall processing of
electric current carrying components prepared by working a
high-strength copper alloy sheet material are taken into account,
it is advantageous that the copper alloy sheet material has an
electrical conductivity of 15% IACS or more. The foregoing
electrical conductivity can be satisfied by the above-described
chemical composition and texture.
[Strength]
[0040] In order to respond to more downsizing and thin-wall
processing of electrical or electronic components using a Cu--Ti
based copper alloy, it is desirable that a 0.2% offset yield
strength in LD is 850 MPa or more. The 0.2% offset yield strength
in LD is controlled to a strength level of more preferably 900 MPa
or more, and still more preferably 950 MPa or more. In addition, a
tensile strength in LD is preferably 900 MPa or more, more
preferably 950 MPa or more, and still more preferably 1,000 MPa or
more. By applying a production condition as described later to an
alloy satisfying the above-described chemical composition, it is
possible to fulfill the above-described strength level at the same
time while keeping the bending workability, fatigue resistance, and
stress relaxation resistance high.
[Bending Workability]
[0041] In order to work the Cu--Ti based copper alloy sheet
material into an electric current carrying component such as a
connector, a lead frame, a relay, and a switch, it is advantageous
that the Cu--Ti based copper alloy sheet material has good bending
workability such that in the 90.degree. W-bending test (width of
test piece: 10 mm) in conformity with JIS H3130, a value of R/t
ratio of a minimum bending radius R to a sheet thickness t at which
cracking does not occur is preferably not more than 2.0, and more
preferably not more than 1.0 in both LD and TD. The bending
workability in LD is a bending workability which is evaluated with
a bending working test piece cut out such that LD is the
longitudinal direction, and the bending axis in that test is TD.
Similarly, the bending workability in TD is a bending workability
which is evaluated with a bending working test piece cut out such
the TD is the longitudinal direction, and the bending axis in that
test is LD.
[Fatigue Resistance]
[0042] In general, the fatigue resistance is evaluated in terms of
the load stress of the test piece and the number of repeated
vibrations until rupture of the test piece occurs (so-called "S--N
curve"). The copper alloy sheet material which is subjective in the
present invention is a copper alloy sheet material having fatigue
resistance such that in the fatigue test in conformity with JIS
22273, in a test piece in which the rolling direction (LD) of the
sheet is the longitudinal direction, a fatigue life at a maximum
load stress of 700 MPa on the test piece surface (the number of
repeated vibrations until rupture of the test piece occurs) is
preferably 500,000 times or more, and more preferably 700,000 times
or more. In the Cu--Ti based copper alloy sheet material, though it
was conventionally considered that it was difficult to make both
the above-described high strength and such excellent fatigue
resistance compatible with each other, it has become possible to
realize this by steps including a precursory treatment as described
later. It is also possible to obtain a copper alloy sheet material
in which the above-described fatigue life is 1,000,000 times or
more.
[Stress Relaxation Resistance]
[0043] As for the stress relaxation resistance, a value of TD is
especially important in an application of an onboard connector or
the like, and therefore, it is desirable to evaluate the stress
relaxation properties in terms of a stress relaxation rate using a
test piece in which the longitudinal direction thereof is TD. In an
evaluation method of stress relaxation properties as described
later, in the case of holding at 200.degree. C. for 1,000 hours,
the stress relaxation rate is preferably not more than 5%, and more
preferably not more than 4%.
<<Production Method>>
[0044] The Cu--Ti based copper alloy sheet material which fulfills
the above-described properties can be produced according to the
following production steps.
[0045] "(Melting and casting).fwdarw.(Hot rolling).fwdarw.(Cold
rolling).fwdarw.(Solution treatment).fwdarw.(Precursory
treatment).fwdarw.(Intermediate cold rolling).fwdarw.(Aging
treatment).fwdarw.(Finish cold rolling).fwdarw.(Low-temperature
annealing)"
[0046] Here, the "precursory treatment" is a heating treatment in a
specified temperature range, which is carried out between the
solution treatment and the aging treatment. This is a heat
treatment in which a so-called precursory modulated structure in
which spinodal decomposition starts to occur slightly before the
generation of a modulated structure (spinodal structure) in the
aging treatment is considered to be formed. Incidentally, while the
description is omitted in the above-described steps, a soaking
treatment (or hot forging) is carried out after the melting and
casting, if desired; facing is carried out after the hot rolling,
if desired; and pickling or grinding, or further degreasing is
carried out after each of the heat treatments, if desired. In
addition, the "intermediate cold rolling" between the solution
treatment and the aging treatment, or the "finish cold rolling" and
the "low-temperature annealing" after the aging treatment may be
omitted as the case may be. The respective steps are hereunder
described.
[Melting and Casting]
[0047] A cast slab may be produced by means of continuous casting,
semi-continuous casting, or the like. In order to prevent oxidation
of Ti from occurring, the production may be carried out in an inert
gas atmosphere or in a vacuum melting furnace.
[Hot Rolling]
[0048] A general hot rolling method of a copper alloy can be
applied. In subjecting the cast slab to hot rolling, by carrying
out an initial rolling pass in a high-temperature region of
700.degree. C. or higher where recrystallization easily occurs, the
casting texture is ruptured, and such is advantageous in
contemplating to homogenize the components and texture. However,
when rolling is carried out at a temperature exceeding 950.degree.
C., there may be the case where cracking occurs in a place where
the melting point decreases, such as a segregated place of the
alloy components. It is necessary to carry out the hot rolling in a
temperature region not exceeding 950.degree. C. In order to surely
carry out the generation of complete recrystallization during the
hot rolling step, it is desirable to carry out the rolling at a
rolling ratio of 60% or more in a temperature region of from
950.degree. C. to 700.degree. C. In order to prevent the formation
and coarsening of the precipitate from occurring, it is effective
to carry out the hot rolling at a final pass temperature of
500.degree. C. or higher. After the hot rolling, it is desirable to
carry out quenching by means of water cooling or the like.
[Cold Rolling]
[0049] In the cold rolling which is carried out before the solution
treatment, it is important to control the rolling ratio to 90% or
more, and it is more preferable to control the rolling ratio to 95%
or more. By subjecting a material worked at such a high rolling
ratio to a solution treatment in the subsequent step, a strain
which is introduced by rolling functions as a nucleus of the
recrystallization, and a crystal grain texture having a uniform
crystal grain diameter is obtained. Incidentally, since an upper
limit of the cold rolling ratio is inevitably restricted by a mill
power or the like, it is not required to be particularly specified.
However, from the viewpoint of preventing edge cracking or the like
from occurring, a good result is easily obtainable at a rolling
ratio of not more than approximately 99%.
[Solution Treatment]
[0050] In the case of a Cu--Ti based copper alloy which is
subjective in the present invention, in particular, it is important
to sufficiently solid-solve a .beta. phase that is a granular
precipitate in a solution treatment. In order to achieve this, it
is effective to increase the temperature to a temperature region of
from 750 to 950.degree. C. and hold it. When the heating
temperature of the solution treatment is too low, the
solid-solution of the coarse granular .beta. phase becomes
insufficient. When the temperature is too high, the crystal grain
becomes coarse. In all of these cases, it becomes difficult to
obtain finally a high-strength material with excellent bending
workability. In addition, in the case where the crystal grain
becomes coarse, even when a precursory treatment as described later
is carried out, a fine .beta. phase hardly precipitates
sufficiently in the grain boundary. In that case, even when aging
is carried out at low temperatures, a coarse precipitate of grain
boundary reaction type is formed. It is desirable to adjust a
heating temperature (maximum ultimate temperature) and a heating
and holding time (in-furnace time) such that an average crystal
grain diameter of the recrystallized grain (a twin boundary is not
considered as the crystal grain boundary) is from 5 to 25 .mu.m.
The average crystal grain diameter of the recrystallized grain is
more preferably from 8 to 20 .mu.m. The recrystallized grain
diameter varies with the cold rolling ratio before the solution
treatment or chemical composition. However, by previously
determining a relationship between the solution treatment heat
pattern and the average crystal grain diameter on each alloy
through an experiment, the holding time of the solution treatment
can be set up. Specifically, for example, in the case of a cold
rolling material having a sheet thickness of from 0.1 to 0.5 mm, an
appropriate condition can be set up within a range where the
furnace temperature is from 750 to 950.degree. C., and preferably
from 780 to 930.degree. C., and the in-furnace time is from 5
seconds to 5 minutes. The average crystal grain diameter after the
solution treatment is reflected in an average crystal grain
diameter of a final product. That is, the average crystal grain
diameter in the final product sheet material is substantially equal
to the average crystal grain diameter after the solution
treatment.
[0051] After completion of the heating process after the solution
treatment, a precursory treatment as a subsequent step can be
carried out utilizing a cooling process from the heating. In
addition, the precursory treatment can also be carried out by after
the solution treatment, once decreasing the temperature to the
vicinity of ordinary temperature, followed by reheating. In that
case, after completion of the heating process subsequent to the
solution treatment, quenching is carried out to at least
200.degree. C. at an average cooling rate of 20.degree. C./sec or
more.
[Precursory Treatment]
[0052] After the solution treatment, the resultant is subjected to
a heat treatment (precursory treatment) of holding at a temperature
ranging from 550 to 730.degree. C. for from 10 to 120 seconds. This
temperature region resides in a temperature range higher than a
temperature region of from 450 to 500.degree. C., in which a
maximum strength is obtained by the formation of a modulated
structure (spinodal structure) in a usual aging treatment of the
Cu--Ti based copper alloy. According to studies made by the present
inventors, when the Cu--Ti based copper alloy after completion of
the solution treatment is held in this temperature region, a fine
granular precipitate of .beta. phase is formed in the crystal grain
boundary and the grain. Then, it has been noted that when the
Cu--Ti based copper alloy of a texture state in which the fine
granular precipitate of .beta. phase is existent is subjected to an
aging treatment, the formation of a precipitate of grain boundary
reaction type is conspicuously suppressed. In addition, it has been
noted that in the Cu--Ti based copper alloy of a texture state held
in a temperature region of from 550 to 730.degree. C. after the
solution treatment, a phenomenon in which a temperature region
where in the subsequent aging treatment, the strength becomes
maximum, namely an appropriate aging treatment temperature range,
shifts towards the low-temperature side occurs. Though reasons for
this have not been sufficiently elucidated yet, it may be
conjectured that by holding at from 550 to 730.degree. C., a
precursory texture structure in which spinodal decomposition starts
to occur slightly is obtained, and the peculiar texture structure
possibly makes it very easy to bring about full-scale formation of
a modulated structure (spinodal structure) from a relatively low
temperature. For that reason, in this specification, the holding at
from 550 to 730.degree. C. which is carried out after the solution
treatment is called "precursory treatment".
[0053] When the holding temperature of the precursory treatment is
too high, the formation amount of the fine granular .beta. phase is
liable to become insufficient. In addition, the crystal grain
easily becomes coarse. When the holding temperature is too low, the
precipitate of grain boundary reaction type precipitates. On the
other hand, when the holding time of the precursory treatment is
too long, the granular .beta. phase becomes coarse, and a lowering
of the strength is easily brought. When the holding time is too
short, the formation amount of the fine granular .beta. phase
becomes small, and an action to strengthen the precipitation by the
.beta. phase cannot be sufficiently enjoyed. After heating and
holding of the precursory treatment, the resultant is quenched to
at least 200.degree. C. at an average cooling rate of 20.degree.
C./sec or more. When the cooling rate to this temperature is slow,
aging occurs in a usual aging treatment temperature region, so that
a merit that the aging temperature can be shifted towards the
low-temperature side cannot be enjoyed.
[0054] The precursory treatment can be carried out utilizing the
cooling process of the solution treatment. In that case, the
treatment may be carried out using a continuous plate feeding line
capable of continuously undergoing the solution treatment and the
precursory treatment.
[0055] Meanwhile, after heating and holding of the solution
treatment, the temperature is decreased to the vicinity of ordinary
temperature, and thereafter, the precursory treatment may also be
carried out. In that case, a heat pattern in which after heating
and holding of the solution treatment, the resultant is quenched to
at least 200.degree. C. at an average cooling rate of 20.degree.
C./sec or more, and the temperature is then increased and held at a
temperature ranging from 550 to 730.degree. C. for from 10 to 120
seconds, followed by quenching to at least 200.degree. C. at an
average cooling rate of 20.degree. C./sec or more, is adopted.
[Intermediate Cold Rolling]
[0056] Prior to the aging treatment, cold rolling can be applied,
if desired. In this specification, the cold rolling at this stage
is called "intermediate cold rolling". The intermediate cold
rolling has an effect for promoting the precipitation during the
aging treatment and is effective for lowering the aging temperature
and shortening the aging time for the purpose of bringing out
necessary properties (e.g., electrical conductivity and hardness).
The rolling ratio of the intermediate cold rolling is required to
be not more than 50%, and the rolling ratio of the intermediate
cold rolling is more preferably not more than 40%. When the rolling
ratio is too high, the bending workability in the TD direction of a
final product is deteriorated. In general, the rolling ratio may be
adjusted within the range of not more than 20%. This cold rolling
step may be omitted.
[Aging Treatment]
[0057] In general, an aging treatment of the Cu--Ti based copper
alloy is frequently carried out at a temperature ranging from 450
to 500.degree. C. at which an action to increase the strength due
to the formation of a modulated structure (spinodal structure)
appears most conspicuously. This temperature range simultaneously
overlaps a temperature region where a precipitate of grain boundary
reaction type is easily formed. For that reason, it was
conventionally difficult to suppress the formation of a precipitate
of grain boundary reaction type in a Cu--Ti high-strength copper
alloy. However, in the case of the Cu--Ti based copper alloy having
gone through the above-described precursory treatment, the
appropriate aging treatment temperature range for the purpose of
obtaining a maximum strength shifts towards the low-temperature
side. As described above, it may be considered that this is
possibly caused due to the fact that a precursory texture structure
in which spinodal decomposition starts to occur slightly is formed
due to the precursory treatment, and full-scale formation of a
modulated structure (spinodal structure) easily occurs from a
relatively low temperature. In consequence, it is possible to carry
out the aging treatment to be adopted herein at a temperature at
which the material temperature reaches from 300 to 430.degree. C.
It is more preferably to carry out the aging treatment at a
temperature ranging from 350 to 400.degree. C. An aging time may
be, for example, set up in the range of from 60 to 900 minutes in a
furnace. In the case of suppressing surface oxidation during the
aging treatment as far as possible, a hydrogen, nitrogen, or argon
atmosphere can be used.
[0058] By combining the above-described precursory treatment with
this aging treatment at low temperatures, the formation of a
precipitate of grain boundary reaction type is conspicuously
suppressed. Examples of reasons for this include the fact that
since a fine granular .beta. phase is already formed in the grain
boundary by the precursory treatment, new precipitation of grain
boundary reaction type hardly occurs; and the fact that the aging
treatment temperature falls outside the temperature region where a
precipitate of grain boundary reaction type is easily formed and is
low. In addition, by going through this aging treatment at this low
temperature, it is possible to increase the strength level to one
equal to or higher than the conventional level. As for the reason
for this, it may be considered that a texture state in which the
amount of the coarse .beta. phase is extremely small before the
aging treatment is present, and a precipitate of grain boundary
reaction type is hardly formed during the aging treatment, and
therefore, the amount of solid-solved Ti in the matrix is kept
high, and as a result, a high action to increase the strength is
possibly exhibited due to a modulated structure on the basis of
fluctuation of the concentration of Ti. In addition, it may be
considered that the existence of a fine granular .beta. phase
formed by the precursory treatment also contributes to
precipitation strengthening.
[Finish Cold Rolling]
[0059] The strength level (in particular, a 0.2% offset yield
strength) can be enhanced by finish cold rolling to be carried out
after the aging treatment. The finish cold rolling can be omitted
in an application in which the requirement of the strength level is
not especially high (for example, the 0.2% offset yield strength is
less than 950 MPa). In the case of carrying out the finish cold
rolling, it is more effective to ensure a rolling ratio of 5% or
more. However, the bending workability in the BW direction (TD) is
easily deteriorated with an increase of the finish cold rolling
ratio. It is necessary to control the rolling ratio of finish cold
rolling to the range of not more than 30%. In general, the finish
cold rolling may be carried out within the range of not more than
20%. A final sheet thickness can be, for example, controlled to
from 0.05 to 1.0 mm. The final sheet thickness is more preferably
from 0.08 to 0.5 mm.
[Low-Temperature Annealing]
[0060] After the finish cold rolling, low-temperature annealing can
be applied for the purposes of decreasing the residual stress of
sheet material or enhancing the bending workability, and enhancing
the stress relaxation resistance due to a decrease of dislocation
on the vacancy or slip plane. It is desirable to set up a heating
temperature such that the material temperature reaches from 150 to
430.degree. C. According to this, it is possible to enhance the
strength, the electrical conductivity, the bending workability, and
the stress relaxation resistance at the same time. When this
heating temperature is too high, the precipitation of grain
boundary reaction type easily occurs. Conversely, when the heating
temperature is too low, the effects for improving the
above-described properties are not sufficiently obtained. It is
desirable to ensure a holding time of 5 seconds or more at the
above-described temperature, and in general, a good result is
obtained in the range within one hour. In the case of omitting the
finish cold rolling, in general, this low-temperature annealing is
omitted, too.
Examples
[0061] Each of copper alloys shown in Table 1 was melted and cast
using a vertical semi-continuous casting machine. The resulting
cast slab was heated at 950.degree. C. and then extracted, and hot
rolling was started. A final pass temperature of the hot rolling
resides between 600.degree. C. and 500.degree. C. A total hot
rolling ratio from the cast slab is about 95%. After the hot
rolling, an oxidized layer as a surface layer was removed (faced)
by means of mechanical grinding, thereby obtaining a rolled sheet
having a thickness of 10 mm. Subsequently, the resulting rolled
sheet was subjected to cold rolling at various rolling ratios of
90% or more and then provided for a solution treatment.
Incidentally, a composition of each of commercially available
materials which were used for comparison is described in Table
1.
[0062] The solution treatment was carried out at a heating
temperature for an in-furnace time shown in Table 2. The in-furnace
time was set to 50 seconds. As for a solution treatment condition,
an appropriate condition under which an average crystal grain
diameter after the solution treatment was from 5 to 25 .mu.m (a
twin boundary is not considered as the crystal grain boundary) was
adopted exclusive of a part of Comparative Examples. As for the
appropriate condition, an optimum temperature was determined
through a preliminary experiment depending upon a composition of
each of alloys of Examples and decided.
[0063] After completion of heating of the solution treatment, a
precursory treatment was carried out utilizing a cooling process
thereof, or cooling to ordinary temperature was carried out by
means of usual water cooling. The precursory treatment utilizing a
cooling process was carried out by a method of dipping a sample
immediately after heating of the solution treatment in a salt bath
adjusted at various temperatures of from 600 to 700.degree. C. and
holding it for a prescribed time, followed by water cooling to the
vicinity of ordinary temperature at a cooling rate of 50.degree.
C./sec or more. In addition, with respect to a part of the samples
which had been cooled to ordinary temperature by means of usual
water cooling, the precursory treatment was carried out by applying
a heat treatment subsequent to the above-described dipping in a
salt bath.
[0064] Subsequently, intermediate cold rolling was carried out
according to need, and an aging treatment was applied at various
temperatures of from 300 to 450.degree. C. An aging time was
adjusted to a time such that the hardness became a peak at each of
the aging temperatures. Thereafter, in a part of the Examples,
finish cold rolling and low-temperature annealing were applied,
thereby preparing test samples. As for the above-described
low-temperature annealing condition, a heating temperature (maximum
ultimate temperature) was set to 420.degree. C., and an in-furnace
time was set to 60 seconds. Incidentally, facing was carried out on
the way according to need, thereby adjusting a sheet thickness of
the test material to 0.15 mm. The production condition is shown in
Table 2.
TABLE-US-00001 TABLE 1 Chemical composition (% by mass) Section No.
Cu Ti Fe Co Ni Others Remark Example 1 Balance 3.25 -- -- -- -- --
according to 2 Balance 4.68 -- -- -- -- -- the present 3 Balance
2.22 -- 0.16 -- -- -- invention 4 Balance 3.21 -- -- -- Zr: 0.10,
P0.03 -- 5 Balance 2.94 -- -- 0.15 B: 0.02 -- 6 Balance 3.26 0.18
-- -- Si: 0.12, Al: 0.08, Zn: 0.46 -- 7 Balance 2.83 -- -- -- Sn:
0.13, Mn: 0.04, V: 0.14 -- 8 Balance 3.06 -- -- 0.12 Cr: 0.12, Mg:
0.06 -- 9 Balance 3.25 -- -- -- -- -- 10 Balance 3.25 -- -- -- --
-- 11 Balance 3.25 -- -- -- -- -- Comparative 21 Balance 3.25 -- --
-- -- -- Example 22 Balance 4.68 -- -- -- -- -- 23 Balance 2.22 --
0.16 -- -- -- 24 Balance 3.21 -- -- -- Zr: 0.10, P: 0.03 -- 25
Balance 2.94 -- -- 0.15 B: 0.02 -- 26 Balance 1.80 -- -- -- Mg:
0.07 -- 27 Balance 5.41 -- 0.13 .0.05 Zn: 0.05 -- 28 Balance 3.24
0.68 -- -- -- -- 29 Balance 3.25 -- -- -- -- -- 30 Balance 3.25 --
-- -- -- -- 31 Balance 3.25 -- -- -- -- -- 32 Balance 3.27 -- -- --
-- Commercially available C1990-1/2H 33 Balance 3.31 -- -- -- --
Commercially available C1990-EH Underlined: Falling outside the
scope of the present invention
TABLE-US-00002 TABLE 2 Production condition Intermediate Finish
cold Solution treatment Precursory treatment cold rolling Aging
treatment rolling Temperature Time Temperature Time Rolling ratio
Temperature Time Rolling ratio Section No. (.degree. C.) (sec)
Cooling process (.degree. C.) (sec) (%) (.degree. C.) (hr) (%)
Example 1 825 50 Carried out precursory treatment 650 50 15 400 5.0
-- according to 2 900 50 Carried out precursory treatment 600 50 5
400 3.5 -- the present 3 785 50 Carried out precursory treatment
700 50 -- 400 7.0 20 invention 4 825 50 Carried out precursory
treatment 650 50 6 400 5.0 -- 5 800 50 Carried out precursory
treatment 675 50 10 400 5.5 10 6 825 50 Carried out precursory
treatment 650 50 -- 400 5.0 10 7 800 50 Carried out precursory
treatment 675 50 -- 400 6.0 15 8 800 50 Carried out precursory
treatment 675 50 10 400 5.5 -- 9 825 50 Water cooling 650 50 15 400
5.0 -- 10 825 50 Carried out precursory treatment 650 50 15 300
12.0 -- 11 825 50 Carried out precursory treatment 650 50 15 430
4.0 -- Comparative 21 825 50 Water cooling -- -- 15 450 5.0 --
Example 22 900 50 Water cooling -- -- 5 450 3.5 -- 23 785 50 Water
cooling -- -- -- 450 7.0 20 24 825 50 Water cooling -- -- 6 450 5.0
25 800 50 Water cooling -- -- 10 450 5.5 10 26 750 50 Carried out
precursory treatment 700 50 -- 400 8.0 15 27 -- -- -- -- -- -- --
-- -- 28 850 50 Carried out precursory treatment 650 50 15 400 5.0
-- 29 960 50 Carried out precursory treatment 650 50 15 400 5.0 --
30 730 50 Carried out precursory treatment 650 50 15 400 3.5 -- 31
825 50 Carried out precursory treatment 650 150 15 400 3.5 -- 32 --
-- -- -- -- -- -- -- -- 33 -- -- -- -- -- -- -- -- --
[0065] In Table 1, No. 32 and No. 33 are concerned with test
materials prepared by obtaining commercially available Cu--Ti based
copper alloys C1990-1/2H and C1990-EH (sheet thickness: 0.15 mm),
respectively. Test pieces were collected from the respective test
materials after the aging treatment or low-temperature annealing,
as obtained in the above-described steps, and the test materials
using a commercially available material (sheet thickness of all of
the materials: 0.15 mm) and examined with respect to an average
crystal grain diameter, a width of a precipitate of grain boundary
reaction type, a density of a granular precipitate having a
diameter of 100 nm or more, an electrical conductivity, a tensile
strength, a 0.2% offset yield strength, fatigue resistance, stress
relaxation resistance, and bending workability.
[0066] The textures and properties were examined in the following
manners.
[Average Crystal Grain Diameter]
[0067] A sheet surface (rolled surface) of the test material was
polished and then subjected to etching, the resulting surface was
observed by an optical microscope, and a grain diameter of 100 or
more crystal grains in a visual field of 300 .mu.m.times.300 .mu.m
was measured by the cutting method of JIS H0501.
[Precipitate of Grain Boundary Reaction Type and Coarse Granular
Precipitate]
[0068] A sheet surface (rolled surface) of the test material was
polished, and the resulting surface was observed by a scanning
electron microscope (SEM, magnification: 3,000 times, observation
field: 42 .mu.m.times.29 .mu.m) in randomly selected five visual
fields.
[0069] A maximum value of a length of the precipitate of grain
boundary reaction type in the rectangular direction to a crystal
grain boundary on which the precipitate of grain boundary reaction
type was formed, the length being measured at a position above the
crystal grain boundary in the five visual fields, was defined as a
maximum width of the precipitate of grain boundary reaction
type.
[0070] A density of the coarse granular precipitate was determined
by dividing the number of granular precipitates having a diameter
of 100 nm or more, as observed in the five visual fields, by a
total area of the visual fields.
[Electrical Conductivity]
[0071] An electrical conductivity of each of the test materials was
measured in conformity with JIS H0505.
[Tensile Strength and 0.2% Offset Yield Strength]
[0072] A tensile test piece (JIS No. 5) of LD was collected from
each of the test materials, subjected to a tensile test at n=3
according to JIS 22241, and measured with respect to a tensile
strength and a 0.2% offset yield strength. The tensile strength and
the 0.2% offset yield strength were determined in terms of an
average value at n=3.
[Bending Workability]
[0073] A bending test piece in which LD is the longitudinal
direction and a bending test piece in which TD is the longitudinal
direction (width of all of the test pieces: 10 mm) were collected
from the sheet material of the test material and subjected to the
90.degree. W-bending test in conformity with JIS H3130. With
respect to the test piece after the test, a surface and a cross
section of the bending-worked part were observed by an optical
microscope with a magnification of 100 times. A minimum bending
radius R at which cracking did not occur was determined, and this
was divided by a sheet thickness t of the test material, thereby
determining an R/t value (MBR/t) of each of LD and TD. The test was
carried out at n=3 in each of LD and TD of each test material, and
a record of the test piece in which the worst result was presented
at n=3 was adopted, thereby expressing an R/t value. Incidentally,
in the case where cracking occurred under a bending condition of
R/t=5.0, a test was not carried out at an R value more than this.
That case is expressed by "ruptured".
[Fatigue Resistance]
[0074] The fatigue test was carried out using a test piece in the
parallel direction to the rolling direction in conformity with JIS
22273. One end of a strip-shaped test piece having a width of 10 mm
was fixed by a fixing tool, and the other end was given sinusoidal
wave vibration via a knife edge, thereby measuring a fatigue life.
A fatigue life at a maximum load stress of 700 MPa on the test
piece surface (the number of repeated vibrations until rupture of
the test piece occurred) was measured. The measurement was carried
out 4 times under the same condition, thereby determining an
average value of the measurement of 4 times.
[Stress Relaxation Properties]
[0075] A bending test piece (width: 10 mm) in which TD was the
longitudinal direction was collected from each of the test
materials and fixed in an arched state such that the surface stress
in a central part in the longitudinal direction of the test piece
was 80% in terms of a 0.2% offset yield strength. The
above-described surface stress is defined according to the
following equation:
Surface stress (MPa)=6Et.delta./L.sub.0.sup.2
wherein
[0076] E: Elastic modulus (MPa)
[0077] t: Thickness of the sample (mm)
[0078] .delta.: Deflection height of the sample (mm)
[0079] A stress relaxation rate was calculated from a bending habit
after holding the test piece in this state in the air at a
temperature of 200.degree. C. for 1,000 hours according to the
following equation:
Stress relaxation rate
(%)=(L.sub.1-L.sub.2)/(L.sub.1-L.sub.0).times.100
wherein
[0080] L.sub.0: Length of the tool, namely a horizontal distance
between the ends of the sample fixed during the test (mm)
[0081] L.sub.1: Length of the sample at the time of starting the
test (mm)
[0082] L.sub.2: Horizontal distance between the ends of the sample
fixed after the test (mm)
[0083] The test sample having this stress relaxation rate of not
more than 5% was evaluated to have high durability as an on-board
connector and decided to be good enough.
[0084] These results are shown in Table 3. LD and TD described in
Table 3 are a direction coincident with the longitudinal direction
of the test piece.
TABLE-US-00003 TABLE 3 Texture Width of Properties Average
precipitate of Density of granular Stress crystal grain precipitate
having Tensile 0.2% offset Bending Fatigue life relaxation grain
boundary a diameter of Electrical strength yield strength
workability (times .times. rate diameter reaction type 100 nm or
more conductivity (MPa) (MPa) (MBR/t) 10,000) (%) Section No.
(.mu.m) (nm) (10.sup.4 number/mm.sup.2) (% IACS) LD LD LD TD LD TD
Example 1 12 <100 2.4 16.8 1065 975 0.0 1.6 214 2.1 according to
2 8 <100 4.5 15.2 1094 1002 0.0 1.0 305 2.4 the present 3 16
<100 1.7 18.6 936 892 0.0 0.7 87 3.2 invention 4 10 <100 3.1
17.2 956 873 0.0 0.0 96 2.8 5 14 <100 1.9 19.6 948 906 0.0 0.7
112 3.4 6 7 <100 2.0 17.2 995 945 0.0 1.6 145 3.1 7 16 <100
2.5 18.2 950 901 0.0 0.0 165 3.3 8 20 <100 1.2 20.6 916 867 0.0
0.0 76 3.6 9 12 <100 2.3 17.0 1070 981 0.0 1.6 222 2.0 10 12
<100 2.1 16.3 1085 992 0.0 1.0 236 2.0 11 12 <100 2.6 17.4
1035 956 0.0 1.0 204 2.3 Comparative 21 12 1850 2.2 12.3 875 764
1.0 2.0 16 4.6 Example 22 8 3680 4.2 11.7 946 843 1.0 3.0 13 4.9 23
16 1240 1.7 12.7 913 849 1.0 1.6 24 4.2 24 10 890 3.0 13.2 782 644
0.0 0.7 0.5 4.4 25 14 1060 1.8 14.1 956 900 1.0 3.0 28 4.5 26 18
<100 1.3 22.5 812 765 0.0 0.7 32 5.2 27 -- -- -- -- -- -- -- --
-- -- 28 4 <100 15.0 17.6 928 845 2.0 3.0 27 5.2 29 65 1150 0.6
14.3 1012 921 1.0 4.0 18 4.6 30 4 <100 26.0 18.4 862 777 1.0 3.0
24 6.2 31 12 <100 18.0 17.4 954 842 1.5 3.0 35 4.4 32 7 1460 3.3
13.1 846 738 1.5 2.0 12 5.8 33 7 1650 3.9 12.4 958 874 2.0 4.0 16
6.2
[0085] As is noted from Table 3, all of the copper alloy sheet
materials according to the present invention have an average
crystal grain diameter of from 5 to 25 .mu.m, a width of a
precipitate of grain boundary reaction type of not more than 500
nm, and a density of a granular precipitate having a diameter of
100 nm or more of not more than 10.sup.5 number/mm.sup.2 and also
have a high strength such that a 0.2% offset yield strength thereof
is 850 MPa or more, good bending workability such that an R/t value
thereof is not more than 2.0 in both LD and TD, and excellent
fatigue resistance such that a fatigue life thereof at a load
stress of 700 MPa is 500,000 times or more. The width of the
precipitates of grain boundary reaction type of the Examples
according to the present invention was specifically less than 100
nm and was on a level of being not substantially perceived.
Furthermore, all of the copper alloy sheet materials according to
the present invention also have excellent stress relaxation
resistance such that the stress relaxation rate of TD which is
important in an application of an onboard connector or the like is
5% or less. In addition, the electrical conductivity of all of the
copper alloy sheet materials according to the present invention is
also improved as compared with C1990 (Nos. 32 and 33) representing
a usual Cu--Ti based copper alloy.
[0086] On the other hand, Comparative Examples Nos. 21 to 25 are
concerned with an example in which the alloys having the same
composition as that in Example Nos. 1 to 5 according to the present
invention were produced by usual steps (those quenched after the
solution treatment), respectively. In all of these Comparative
Examples, the formation of a precipitate of grain boundary reaction
type is not suppressed, and the strength, bending workability,
fatigue resistance, stress relaxation resistance, electrical
conductivity, and the like are generally inferior to those in the
Examples according to the present invention.
[0087] Comparative Examples Nos. 26 to 28 are concerned with an
example in which good properties were not obtained in view of the
fact that the chemical composition falls outside the scope of the
present invention. No. 26 is low in the strength level and inferior
in the fatigue resistance because of an excessively low content of
Ti. No. 27 could not take an appropriate solution treatment
condition because of an excessively high content of Ti, so that
cracking occurred on the way of production, and a sheet material
capable of being evaluated could not be prepared. No. 28 was
substantially free from the precipitation of grain boundary
reaction type because Fe was added for the purpose of suppressing
the precipitation of grain boundary reaction type; however, Fe and
Ti formed a coarse intermetallic compound (granular precipitate) in
view of the fact that the addition amount of Fe was in excess, and
all of the strength, bending workability, fatigue resistance, and
stress relaxation resistance were deteriorated.
[0088] Comparative Examples Nos. 29 to 31 are concerned with an
example in which good properties were not obtained in view of the
fact that with respect to the alloy having the same composition as
that in Example No. 1 according to the present invention, the
heating and holding condition of the solution treatment or the
precursory treatment condition falls outside the scope of the
present invention. In No. 29, the crystal grain was coarsened
because of an excessively high heating temperature of the solution
treatment relative to the holding time of 50 seconds, and
nevertheless the precursory treatment was applied during the
subsequent cooling, the progress of precipitation of grain boundary
reaction type was not sufficiently suppressed during the aging
treatment. As a result, good fatigue resistance was not obtained.
In addition, the bending workability was inferior due to coarsening
of the crystal grain. In No. 30, since the solution treatment
temperature was conversely too low as 730.degree. C., a large
amount of the granular precipitate having a diameter of 100 nm or
more remained (in a solid-unsolved state). In that case, though the
precipitation of grain boundary reaction type could be suppressed
during the aging treatment, bad results were brought in all of the
strength, fatigue resistance, bending workability, and stress
relaxation resistance. In No. 31, since the holding time of the
precursory treatment was too long, the granular precipitate was
excessively formed. As a result, though the precipitation of grain
boundary reaction type could be suppressed during the aging
treatment, the strength, fatigue resistance, and bending
workability were inferior.
[0089] Comparative Examples Nos. 32 and 33 are commercially
available products of C1990-1/2H and C1990-EH representing the
Cu--Ti based copper alloy. In all of them, a precipitate of grain
boundary reaction type having a width exceeding 500 nm is formed,
and as compared with Example No. 1 according to the present
invention having substantially the same composition, all of the
strength, fatigue resistance, bending workability, stress
relaxation resistance, and electrical conductivity are
inferior.
[0090] An SEM photograph of a cross section perpendicular to the
sheet thickness direction with respect to the test material of
Comparative Example No. 21 which was produced in usual steps is
illustrated in FIG. 2. In addition, an SEM photograph similar to
that in FIG. 2 with respect to the test material of Example No. 1
according to the present invention using an alloy having the same
composition as that in FIG. 2 is illustrated in FIG. 3. In FIG. 2
(Comparative Example), a large number of precipitates of grain
boundary reaction type having a width largely exceeding 500 nm are
observed. On the other hand, the presence of a precipitate of grain
boundary reaction type is not confirmed in FIG. 1 (Example
according to the present invention).
* * * * *